JP7196667B2 - Manufacturing method of sintered body for rare earth magnet - Google Patents

Description

The present invention relates to a method for producing a sintered body for rare earth magnets.

Permanent magnets are used in various motors such as automobile parts, industrial machinery, and home appliances.

Nd--Fe--B system magnets can be cited as typical high-performance magnets. Nd--Fe--B magnets are mainly used in drive motors for electric vehicles (EV, HV, PHV, etc.) and hybrid vehicles. As the need for motors with higher efficiency and smaller size increases, the development of permanent magnets with higher magnetic properties is expected.

An RT ₁₂ compound having a ThMn ₁₂ type crystal structure or a structure similar thereto is attracting attention as one of the candidates for main phase system alloys for permanent magnets that surpasses the magnetic properties of Nd--Fe--B magnets. The RT ₁₂ compound is R ₂ T ₁₄ B (R is at least one rare earth element, T is at least one iron group transition metal element containing at least Fe), which is a compound that constitutes the main phase of an Nd—Fe—B magnet. ) Higher magnetic properties are expected due to higher concentrations of iron group transition metals. Hereinafter, a phase composed of an RT ₁₂ compound having a ThMn ₁₂ -type crystal structure or a similar structure may be referred to as a 1-12 phase.

Patent Document 1 discloses a rare earth permanent magnet in which a portion of Fe, which is a T element, is partially replaced with Ti, which is a structure stabilizing element, to improve thermal stability in exchange for high magnetization. there is

In Patent Document 2, by partially substituting the R element of the RFe ₁₂ -based compound with a substituting element such as Zr or Hf, the amount of the substituting element such as Ti substituting the transition metal element is reduced to increase the saturation magnetization. Rare earth permanent magnets are disclosed which are stabilized while preserving the ThMn type ₁₂ crystal structure.

Further, Patent Document 3 discloses an R'--Fe--Co system ferromagnetic alloy in which Y or Gd is selected as part of the R element of RFe ₁₂ , and this R'--Fe--Co system ferromagnetic It is described that the alloy exhibits high magnetic properties due to having a ThMn ₁₂ -type crystal structure produced by an ultraquenching method.

Further, Patent Document 4 describes that by adding Cu, a non-magnetic and low melting point 1-4 composition (SmCu ₄ phase) phase is generated, and sintering and high coercive force are possible. .

Further, in Patent Document 5, at least one of Sm ₅ Fe ₁₇ -based phase, SmCo ₅ -based phase, Sm ₂ O ₃ -based phase, and Sm ₇ Cu ₃ -based phase is added as a secondary phase to a ThMn ₁₂ -type main phase. It is described that the coercive force can be increased by including the element.

Further, Patent Document 6 describes that the addition of Cu generates a liquid phase and enables the formation of a dense bulk body.

In addition, in Patent Document 7, by producing a ferromagnetic alloy whose main phase is a ThMn ₁₂ type phase containing Y by a strip casting method, the main phase composition is less uniform and the main phase ratio is high. is obtained.

Further, Patent Document 8 describes that a magnetic material having a Y-containing ThMn ₁₂ -type phase as a main phase can provide high saturation magnetization and anisotropic magnetic field.

JP-A-64-76703 JP-A-4-322406 JP 2015-156436 A Japanese Patent Application Laid-Open No. 2001-189206 Japanese Unexamined Patent Application Publication No. 2017-112300 WO2016/162990 JP 2018-103211 A JP 2018-125512 A

As one of the conditions for a sintered body to be used for a high-performance magnet, it is necessary to have a structure with few heterogeneous phases that adversely affect magnetic properties. When a soft magnetic phase typified by the bcc-Fe phase exists in the sintered body, the soft magnetic phase becomes a starting point for magnetization reversal, and magnetization reversal easily progresses. Magnetic properties are significantly degraded. Therefore, a sintered body in which such a heterogeneous phase of soft magnetism does not exist as much as possible is required. Moreover, it is necessary to orient the sintered body in order to increase the residual magnetic flux density (B _r ). For that purpose, it is required that the amount of the main phase is large in the state of fine powder, and that the main phase is a single-crystal-like particle. It is required to be sufficiently coarse with respect to the size of

The rare earth permanent magnet described in Patent Document 1 has improved thermal stability due to element substitution of Fe with Ti, but since the amount of Fe substituted with Ti is large, the magnetization is correspondingly reduced, and sufficient magnetic properties are obtained. I can't get it.

On the other hand, in the rare earth permanent magnet described in Patent Document 2, the _ThMn12 structure is stabilized by substituting a transition metal element with Ti or the like, but the effect is not necessarily sufficient.

The R'--Fe--Co ferromagnetic alloy described in Patent Document 3 does not replace the Fe element with the structure stabilizing element M (such as Ti), so it exhibits high magnetization, large magnetic anisotropy, and a high Curie temperature. However, due to the non-equilibrium phase, the main phase compound may decompose in a high temperature densification process such as sintering.

The rare earth magnet described in Patent Document 4 may not have high magnetic physical properties due to the large amount of Ti added.

In the rare earth magnet described in Patent Document 5, when the rare earth-rich secondary phase Sm ₇ Cu ₃ is used, a phase richer in rare earth than the main phase may be generated due to the reaction between the main phase and Sm ₇ Cu ₃ during heat treatment. Concerned.

In the rare earth magnet described in Patent Document 6, since the Fe element is not replaced with the structure stabilizing element M, high magnetization, large magnetic anisotropy, and a high Curie temperature can be obtained, and the density as a bulk body is high. Since it is a non-equilibrium phase, the main phase compound may decompose in a high temperature process such as sintering at 1000° C. or higher.

Since the composition of the ferromagnetic alloy described in Patent Document 7 and the magnet material described in Patent Document 8 does not take into account the influence of oxygen that is inevitably mixed in the production process of the sintered body, oxygen is preferred as a rare earth element. It is feared that a soft magnetic phase such as a bcc-Fe phase may be generated due to the decomposition of the main phase.

Embodiments of the present disclosure provide a method for manufacturing a sintered body for a rare earth magnet that has less heterogeneous phases that adversely affect magnetic properties and that can be oriented.

In an exemplary embodiment of the method for producing a sintered body for rare earth magnets of the present disclosure, the overall composition is represented by the following compositional formula (1),
R1 _1-x1 R2 _x1 (Fe _1-y1 Co _y1 ) _w1-z1 Ti _z1 Cu _α1 O _β1 (1),
R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and must contain Sm, Sm is 50 mol% or more of the total of R2, and x ₁ , y ₁ , z ₁ , w ₁ , α ₁ , and β ₁ are 0.3 ≤ x ₁ ≤ 0.9, 0 ≤ y ₁ ≤ 0, respectively .4, 0.38≦z ₁ ≦0.70, 7≦w ₁ ≦12, 0≦α ₁ ≦0.70, 0.02≦β ₁ ≦0.5, and 0≦1−x ₁ −2z A method for producing a sintered body for a rare earth magnet having a phase having a ThMn type ₁₂ crystal structure as a main phase and satisfying ₁ /3-0.092α ₁ -8β ₁ /15≦0.05, comprising: is cooled, and the overall composition is represented by the following compositional formula (2),
R3 _1-x2 R4 _x2 (Fe _1-y2 Co _y2 ) _w2-z2 Ti _z2 Cu _α2 O _β2 (2),
R3 is Y or Y and Gd, Y is 50 mol% or more of the total of R3, R4 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and must contain Sm, Sm is 50 mol% or more of the total of R4, and x ₂ , y ₂ , z ₂ , w _{2 ,} α ₂ and β ₂ are respectively 0.3≦x ₂ ≦0.9, 0≦y ₂ ≦0.4 , 0.38≦z ₂ ≦0.70, 9.4≦w ₂ ≦12.0, 0.44≦α ₂ ≦0.70, and β ₂ ≦0.5. The process of making and cooling the molten metal of the raw material, the overall composition is represented by the following compositional formula (3),
R5 _1-x3 R6 _x3 T3 _w3 Cu _α3 O _β3 (3),
R5 is Y or Y and Gd, Y is 50 mol% or more of the total of R5, R6 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always containing Sm, Sm is 50 mol% or more of the total of R6, T3 is at least one selected from the group consisting of Fe, Co, and Ti, and x ₃ , w _{3 ,} α ₃ and β ₃ are each 0≦x ₃ < 0.5, 0≦w ₃ ≦3, 0≦α ₃ ≦2, and β ₃ ≦0.5; a step of mixing at a mixing ratio of 10% or less to obtain a mixed fine powder of alloy M and alloy L; a step of producing a compact of the mixed fine powder; a step of heat-treating for minutes to 50 hours to obtain a sintered body.

In one embodiment, in the powder X-ray diffraction pattern of the sintered body, the maximum intensity of the peak due to the 002 reflection of the phase having the ThMn ₁₂ type crystal structure is I _ThMn12 , bcc-(Fe, Co, Ti) phase where _{Ibcc- ( Fe,Co,Ti)} is the maximum intensity of the peak due to the 011 reflection of .

In one embodiment, in the powder X-ray diffraction pattern of the sintered body, the maximum intensity of the peak due to the 002 reflection of the phase having the ThMn ₁₂ type crystal structure is defined as I _ThMn12 , the phase having the Th ₂ Ni ₁₇ type crystal structure where I _Th2Ni17 is the maximum intensity of the peak due to the 023 reflection of , I _Th2Ni17 /I _ThMn12 ≦0.7 is satisfied.

According to the embodiments of the present invention, it is possible to provide a method for producing a sintered body for a rare earth magnet that has less heterogeneous phases that adversely affect magnetic properties and that can be oriented.

Sample no. FIG. 4 is a diagram showing backscattered electron (BSE) images of cross sections of alloys after heat treatment in M1 to M8. Sample no. 1 is a diagram showing powder X-ray diffraction measurement results in 1 to 4. FIG. Sample no. 1 to 11, the relative intensity of the bcc-(Fe, Co, Ti) phase obtained from the powder X-ray diffraction measurement results for the value of 1-x ₁ -2z ₁ /3-0.092α ₁ -8β ₁ /15 It is a figure which shows. Sample no. 1 is a diagram showing the relative intensity of the 2-17 phase obtained from powder X-ray diffraction measurement results with respect to the values of 1-x ₁ -2z ₁ /3-0.092α ₁ -8β ₁ /15 in 1 to 11. FIG. Sample no. 1 is a diagram showing BSE images of sample cross sections in 1 to 11. FIG.

As a result of intensive research by the present inventors, it was found that by making the final composition of the sintered body satisfy the composition formula (1) described later, the bcc-(Fe, Co, Ti) phase and 2 It was found that the formation of -17 phase and the like can be suppressed. However, if an attempt is made to produce such a sintered body from only one kind of alloy, a large amount of R-rich phase such as 2-17 phase is produced at the stage of the alloy, which is richer in R than 1-12 phase. In that case, the ratio of the 1-12 phase in the alloy is low, and the 1-12 phase is not sufficiently coarse with respect to the pulverization size, so that it is not well oriented when compacted in a magnetic field. The present inventors separately produced a raw material alloy M that satisfies the composition formula (2) described later and a raw material alloy L that satisfied the composition formula (3) described later, and mixed them in the middle to obtain the orientation. In addition, the sintered body also satisfies the composition formula (1) to suppress phases that adversely affect the magnetic properties.

[Composition of sintered body for rare earth magnet]
The overall composition of the sintered body for a rare earth magnet of the present disclosure is represented by the following compositional formula (1).
R1 _1-x1 R2 _x1 (Fe _1-y1 Co _y1 ) _w1-z1 Ti _z1 Cu _α1 O _β1 (1)

Here, R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and Sm Sm must be included, and Sm is 50 mol % or more of the entire R2. R1 is preferably Y only (excluding unavoidable impurities), and R2 is preferably Sm only (excluding unavoidable impurities). In addition, x ₁ , y ₁ , z ₁ , w ₁ , α ₁ and β ₁ are respectively 0.3≦x ₁ ≦0.9, 0≦y ₁ ≦0.4, 0.38≦z ₁ ≦0 .70, 7≦w ₁ ≦12, 0.4≦α ₁ ≦0.7 and 0.02≦β ₁ ≦0.5, and the relational expression 0≦1−x ₁ −2z ₁ /3− It satisfies 0.092α ₁ -8β ₁ /15≦0.05.

As a result of intensive research by the present inventors, by setting the composition range of the sintered body as shown in the above formula (1), the bcc-(Fe, Co, Ti) phase, which adversely affects the magnetic properties, and It was found that the formation of 2-17 phase and the like can be suppressed.

[Regarding the reason for limiting the composition of the sintered body, etc.]
(Types of R1 and R2)
R1 is Y or Y and Gd, and Y is 50 mol % or more of the total of R1. When R1 is another element, a stable phase other than the 1-12 phase may be produced. For example, when R1 is Zr, a Th ₆ Mn ₂₃ type phase is produced, and a large amount of bcc-(Fe, Co, Ti) phase is also produced, making it impossible to obtain a desired sintered body. Since Gd is expensive, it is preferable that R1 is only Y. In addition, from the magnetic physical property value and phase stability of the 1-12 phase, R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is the total of R2 50 mol % or more. From the viewpoint of magnetic physical properties, R2 is more preferably Sm only.

(Content of R2)
The range of x ₁ (R2 substitution amount x ₁ ), which indicates the atomic number ratio of the content of R2 to the total amount of R1 and R2, is 0.3≦x ₁ ≦0.9. If x ₁ − is less than 0.3, the magnetic anisotropy of the 1-12 phase is lowered, which is not preferable. On the other hand, if x1 is larger than 0.9, the stability of the _1-12 phase is lowered, and the bcc-(Fe, Co, Ti) phase and the 2-17 phase may be generated, which is not preferable.

(ratio of Fe and Co)
The range of y ₁ (Co substitution amount y ₁ ), which indicates the atomic number ratio of Co to the sum of Fe and Co, is 0≦y ₁ ≦0.4. In order to avoid the risk of lowering the Curie temperature of the _1-12 phase, y1 is more preferably 0.05 or more. Also, if y1 is larger _than 0.4, the volume magnetization and magnetic anisotropic magnetic field of the 1-12 phase are lowered, which is not preferable.

(Content of Ti)
The range of z ₁ (Ti content z ₁ ), which indicates the atomic number ratio of the Ti content to the total amount of R1 and R2, is 0.38≦z ₁ ≦0.70. If z1 is _less than 0.38, the 2-17 phase and the bcc-(Fe, Co, Ti) phase are stably generated during sintering, which is not preferable. Also, when z1 is larger than 0.70, the magnetic properties of the _1-12 phase are lowered, which is not preferable. In order to obtain higher magnetic properties, especially _Js , the smaller the Ti content, the better. Specifically, it is more preferable that the range of z ₁ is 0.38≦z ₁ ≦0.60. Note that 50 mol% or less of Ti is an element that stabilizes the 1-12 phase structure, such as tungsten (W), vanadium (V), niobium (Nb), tantalum (Ta), molybdenum (Mo), and silicon (Si). may be replaced.

(Cu content)
_The range of α1, which indicates the atomic number ratio of the Cu content to the total amount of _R1 and R2, is 0.4≦α1≦0.7. If α1 is larger _than 0.7, the ratio of the R—Cu phase, which is a secondary phase, increases, the ratio of the main phase decreases, and the magnetization of the sintered body as a whole decreases, which is not preferable. On the other hand, when α1 is _less than 0.4, the amount of liquid phase during heat treatment is reduced, which makes it difficult to reduce heterogeneous phases during solution treatment and densification during sintering, which is not preferable.

(Total amount of Fe, Co and Ti)
The range of w1, which indicates the atomic number ratio of the total amount of Fe, Co, and Ti to the total amount of _R1 and R2, is _7≤w1≤12 . When w1 is larger _than 12, the bcc-(Fe, Co, Ti) phase is produced significantly, which is not preferable. If w1 is _less than 7, a phase such as the 2-17 phase, which has a higher rare earth element content than the 1-12 phase and has an adverse effect on the magnetic properties, is not preferable.

(Oxygen content)
β1, which indicates the atomic number ratio of the oxygen content to the _total amount of _R1 and R2, is suitably in the range of 0.02≦β1≦0.5. If β1 is _less than 0.02, the fine powder before sintering tends to ignite, making handling difficult, which is not preferable. On the other hand, if β1 is larger _than 0.5, the ratio of the oxide phase in the sintered body increases, the ratio of the 1-12 phase decreases, and the magnetization of the magnet as a whole decreases, which is not preferable.

(Relationship between the amount of oxygen and the amount of other elements)
x ₁ , z ₁ , α ₁ , and β ₁ satisfy the relationship 0≦1−x ₁ −2z ₁ /3−0.092α ₁ −8β ₁ /15≦0.05. Since the sintered body generally uses fine powder, it usually has a higher oxygen content than the raw material alloy. Therefore, even in alloys with few heterogeneous phases at the raw material alloy stage, oxygen reacts with rare earth elements in the main phase and grain boundary phase (liquid phase during sintering) during pulverization and sintering to form oxide phases. As the 1-12 phase decomposes, the bcc-(Fe, Co, Ti) phase may be generated. As a result of intensive research, the authors found out that Y contained in R1 is particularly easily oxidized and how R1 is distributed in each phase. From the value of z ₁ , the amount of R1 consumed as the 1-12 phase is 2z ₁ /3, and from the value of α ₁ , the amount of R1 consumed as the R-Cu phase is 0.092α ₁ , β ₁ We describe the R1 consumed as the R-oxide phase from the values as 8β ₁ /15 respectively, and the actual amount of R1 1−x ₁ minus the amount of R1 calculated from z ₁ , α ₁ , β ₁ above This formula defines the lower limit. As 1-x ₁ -2z ₁ /3-0.092α ₁ -8β ₁ /15 becomes smaller on the negative side, R1 necessary for the formation of 1-12 phase, R-Cu phase and R oxide phase is insufficient. In particular, when it is less than 0, a large amount of bcc-(Fe, Co, Ti) phase is generated, which is not preferable. Conversely, the larger the value on the plus side, the more R1 becomes excessive, and especially when it is larger than 0.05, it is not preferable because a large amount of phases such as 2-17 phases having a higher rare earth content than 1-12 phases are generated. . Moreover, powder X-ray diffraction measurement can be mentioned as a simple method for examining the extent to which the bcc-(Fe, Co, Ti) phase and 2-17 phase are present in the sintered body. Among the peaks of each phase, the phase having a ThMn ₁₂ type crystal structure (1-12 phase) has 002 reflection, the bcc-(Fe, Co, Ti) phase has 011 reflection, and the phase having Th ₂ Ni ₁₇ type crystal structure ( 2-17 phase), the peak due to 023 reflection is less affected by other phases, and the peak intensity is high. Therefore, the maximum intensity of the peak due to the 002 reflection of the 1-12 phase is I _ThMn12 , and the maximum intensity of the peak due to the 011 reflection of the bcc-(Fe, Co, Ti) phase is I _{bcc-(Fe, Co, Ti )} , the relative intensities of the peaks of the _bcc- (Fe, Co, Ti) phase and the 2-17 phase are I _{bcc- (Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 . It is desirable that the bcc-(Fe, Co, Ti) phase and the 2-17 phase in the sintered body are as small as possible, and the relative strength of the bcc-(Fe, Co, Ti) phase I _{bcc-(Fe, Co, Ti)} /I _ThMn12 is preferably 0.75 or less. Also, the relative intensity I _Th2Ni17 /I _ThMn12 of the 2-17 phase is preferably 0.7 or less.

[Composition of Alloy M for Rare Earth Magnet]
The overall composition of the rare earth magnet alloy M of the present disclosure (hereinafter sometimes simply referred to as “alloy M” or “raw alloy M”) is represented by the following compositional formula (2).
R3 _1-x2 R4 _x2 (Fe _1-y2 Co _y2 ) _w2-z2 Ti _z2 Cu _α2 O _β2 (2)

Here, R3 is Y or Y and Gd, Y is 50 mol% or more of the total of R3, R4 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and Sm It must be included, and Sm is 50 mol % or more of the entire R4. R3 is preferably Y only (excluding unavoidable impurities), and R4 is preferably Sm only (excluding unavoidable impurities). In addition, x ₂ , y ₂ , z ₂ , w ₂ , α ₂ and β ₂ are 0.3≦x ₂ ≦0.9, 0≦y ₂ ≦0.4, 0.38≦z ₂ ≦0, respectively. .70, 9.4≦w ₂ ≦12.0, 0.44≦α ₂ ≦0.70 and β ₂ ≦0.5.

As a result of intensive research by the present inventors, by setting the composition range of the rare earth magnet alloy M as shown in the above formula (1), the alloy after heat treatment has a crystal grain size of 1-12 μm. It was found that a structure consisting of a phase and an R--Cu grain boundary phase was formed, and that the subsequent fine pulverization gave a fine powder in which the 1-12 phase was like a single crystal.

[Reasons for limiting the composition of alloy M]
(Types of R3 and R4)
R3 is Y or Y and Gd, and Y is 50 mol % or more of the total of R3. When R3 is another element, a stable phase other than the 1-12 phase may be generated. Since Gd is expensive, it is preferable that R3 is only Y. In addition, from the magnetic property value and phase stability of the 1-12 phase, R4 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is the total of R2 50 mol % or more. From the viewpoint of magnetic physical properties, R4 is more preferably Sm only.

(Content of R4)
The range of x ₂ (R4 substitution amount x ₂ ), which indicates the atomic number ratio of the content of R4 to the total amount of R3 and R4, is 0.3≦x ₂ ≦0.9. If x2 is less _than 0.3, the magnetic anisotropy of the 1-12 phase is lowered, which is not preferable. On the other hand, if x2 is larger _than 0.9, the stability of the 1-12 phase is lowered, and the bcc-(Fe, Co, Ti) phase and the 2-17 phase may be generated, which is not preferable.

(ratio of Fe and Co)
The range of y ₂ (Co substitution amount y ₂ ), which indicates the atomic number ratio of Co to the sum of Fe and Co, is 0≦y ₂ ≦0.4. In order to avoid the risk of lowering the Curie temperature of the 1-12 phase, y2 is _more preferably 0.05 or more. Also, if y2 is larger _than 0.4, the volume magnetization and magnetic anisotropic magnetic field of the 1-12 phase are lowered, which is not preferable.

(Content of Ti)
The range of z ₂ (Ti content z ₂ ) indicating the atomic number ratio of the Ti content to the total amount of R3 and R4 is 0.38≦z ₂ ≦0.70. If z2 is less than 0.38, the _2-17 phase and the bcc-(Fe, Co, Ti) phase are stably generated during sintering, which is not preferable. Moreover, if z2 is larger _than 0.70, the magnetic properties of the 1-12 phase are lowered, which is not preferable. In order to obtain higher magnetic properties, especially _Js , the smaller the Ti content, the better. Specifically, it is _more preferable that the range of z2 is 0.38≦ _z2 ≦0.60. Note that 50 mol % or less of Ti may be replaced with elements such as W, V, Nb, Ta, Mo, and Si that stabilize the structure of the 1-12 phase.

(Cu content)
_The range of α2, which indicates the atomic number ratio of the Cu content to the _total amount of R3 and R4, is 0.44≦α2≦0.70. If α2 is larger _than 0.70, the ratio of the R—Cu phase, which is the secondary phase, increases, the ratio of the main phase decreases, and the magnetization of the sintered body as a whole decreases, which is not preferable. Also, if α2 is less _than 0.44, the reduction of heterogeneous phases during heat treatment of the alloy and the coarsening of the 1-12 phase grain size become difficult to progress, which is not preferable.

(Total amount of Fe, Co and Ti)
The range of w2, which indicates the atomic number ratio of the total amount of Fe, Co, and Ti to the _total amount of R3 and R4, is _{9.4≤w2≤12.0} . If w2 is larger _than 12.0, the bcc-(Fe, Co, Ti) phase is produced significantly, which is not preferable. Also, when w2 is less than 9.4, a phase such as the _2-17 phase having a higher rare earth content than the 1-12 phase is generated, and the 1-12 phase is less likely to coarsen during the alloy heat treatment. I don't like it.

(Oxygen content)
β ₂ , which indicates the atomic number ratio of the oxygen content to the total amount of R3 and R4, is suitably in the range of β ₂ ≦0.5. If β2 is _more than 0.5, the ratio of the oxide phase in the sintered body increases, the ratio of the 1-12 phase decreases, and the magnetization of the magnet as a whole decreases, which is not preferable.

[Composition of alloy L for rare earth magnet]
The overall composition of the rare earth magnet alloy L of the present disclosure (hereinafter sometimes simply referred to as “alloy L” or “raw material alloy L”) is represented by the following compositional formula (3).
R5 _1-x3 R6 _x3 T3 _w3 Cu _α3 O _β3 (3)

Here, R5 is Y or Y and Gd, Y is 50 mol% or more of the total of R5, R6 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and Sm It must be included, and Sm is 50 mol % or more of the entire R6. R5 is preferably Y only (excluding unavoidable impurities), and R6 is preferably Sm only (excluding unavoidable impurities). T3 is at least one selected from the group consisting of Fe, Co and Ti. Also, x ₃ , w ₃ , α ₃ and β ₃ satisfy 0≦x ₃ <0.5, 0≦w ₃ ≦3, 0≦α ₃ ≦2 and β ₃ ≦0.5 respectively.

[Reasons for limiting the composition of alloy L]
(Types of R5 and R6)
R5 is Y or Y and Gd, and Y is 50 mol % of all R5. Since Gd is expensive, it is preferable that R5 is only Y. Also, during sintering, a portion of the rare earth elements contained in alloy L are used to form the 1-12 phase by diffusion. From the magnetic physical property value and phase stability of the 1-12 phase, R6 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, Sm is 50 mol% of the total R6 That's it. From the viewpoint of magnetic physical properties, R6 is more preferably Sm only.

(Content of R6)
The range of x ₃ (R6 substitution amount x ₃ ), which indicates the atomic number ratio of the content of R6 to the total amount of R5 and R6, is 0≦x ₃ <0.5. During sintering, the reaction of the rare earth element to form an oxide proceeds. At that time, R5 is oxidized preferentially over R6, so it is necessary to increase the atomic number ratio of the content of R5. If _x3 is 0.5 or more, oxidation of the rare earth element in the 1-12 phase proceeds during sintering, and the 1-12 phase is decomposed to form the bcc-(Fe, Co, Ti) phase. I don't like it.

(Type and content of T3)
T3 is at least one element selected from the group consisting of Fe, Co, and Ti, which are elements constituting the 1-12 phase. The range of w ₃ (T3 content w ₃ ), which indicates the atomic number ratio of the content of T3 to the total amount of R5 and R6, is 0≦w ₃ ≦3. T3 is not an essential element of alloy L. Also, if w3 is larger than ₃ , the melting point of the alloy L becomes high, and the entire amount does not become a liquid phase during sintering, making it difficult for diffusion and reaction to proceed, which is not preferable. As T3, an element such as W, V, Nb, Ta, Mo, and Si that stabilizes the 1-12 phase structure may be added instead of Ti.

(Cu content)
The range of α3 _, which indicates the atomic number ratio of the Cu content to the total amount of R5 and R6, is 0≦α3≦ ₂ . Cu is not an essential element of Alloy L. If α3 is larger than ₂ , the ratio of the R—Cu phase, which is the secondary phase, increases, the ratio of the main phase decreases, and the magnetization of the sintered body as a whole decreases, which is not preferable.

(Oxygen content)
β3 _, which indicates the atomic number ratio of the oxygen content to the total amount of R5 and R6, is suitably in the range of _β3 ≦0.5. If β3 is more than 0.5 _, the ratio of the oxide phase in the sintered body increases, the ratio of the 1-12 phase decreases, and the magnetization of the magnet as a whole decreases, which is not preferable.

[Method for producing sintered body for rare earth magnet]
<Step A> Step of Producing Alloy M As a method for producing alloy M, which is a raw material for a sintered body for a rare earth magnet, there are known methods such as die casting, centrifugal casting, strip casting, and liquid ultra-quenching. can be adopted. These methods involve making a molten alloy and then cooling and solidifying the molten alloy. It is desirable to minimize the formation of coarse bcc-(Fe, Co, Ti) phases and 2-17 phases during solidification of the molten alloy. Coarse bcc- It is possible to suppress the generation of (Fe, Co, Ti) phases and 2-17 phases. When the cooling rate during solidification is low, the grain size of the precipitated heterogeneous phase becomes large. When the grain size of the hetero-phase contained in the alloy becomes large, it becomes difficult to eliminate the hetero-phase during heat treatment or sintering of the alloy. Heat treatment may be performed for the purpose of reducing heterogeneous phases generated during the solidification process of alloy M and coarsening the 1-12 phase. Depending on the composition of the alloy, the R—Cu phase melting point is 850-900°C. Therefore, the heat treatment temperature is preferably 900° C. or higher and 1250° C. or lower, more preferably 1000° C. or higher and 1150° C. or lower. Moreover, although the heat treatment time depends on the heat treatment temperature, it is desirable that the heat treatment time is 5 minutes or more and 50 hours or less. If the time is too short, there will not be enough reaction to eliminate the heterophases. If the time is too long, evaporation and oxidation of the rare earth elements will occur, and the operational efficiency will be poor. It is preferable that the 2-17 phase present in the alloy M is as small as possible, and the 2-17 phase is preferably 10% or less in terms of area ratio of the cross section of the alloy. If the 2-17 phase exceeds 10%, the grain size of the 1-12 phase becomes fine, which is not preferable. Additionally, prior to the grinding step, the alloy may be heat treated in hydrogen to introduce cracks. When containing Cu, the R—Cu phase in the alloy can absorb and release hydrogen. With this alloy, for example, hydrogen absorption occurs at temperatures between 250°C and 400°C, and hydrogen release occurs between 540°C and 660°C. Therefore, after the alloy is heated to 400° C. or higher in hydrogen to absorb hydrogen, the atmosphere can be switched to a vacuum atmosphere to sufficiently release hydrogen. In that case, the temperature for switching to the vacuum atmosphere is 700° C. or less. Thus, the subphase contained in the present alloy undergoes hydrogen absorption and desorption at a temperature of at least 700° C. or less. If the alloy is exposed to a hydrogen atmosphere at a temperature above 700° C., decomposition of the main phase may occur due to a hydrogenation-disproportionation reaction. By absorbing and releasing hydrogen, the rare earth-rich phase (subphase) expands and contracts in volume, and cracks occur between the main phase grains and the subphase. This increases the pulverization efficiency in the pulverization process.

<Step B> Step of Producing Alloy L Methods for producing alloy L, which is the raw material of the sintered body for rare earth magnets, include die casting, centrifugal casting, strip casting, liquid ultra-quenching, and atomizing. A known method can be adopted. These methods involve making a molten alloy and then cooling and solidifying the molten alloy. It is desirable to prevent segregation during solidification so that the alloy has a uniform composition throughout. A method with relatively high cooling speed, such as strip casting method or liquid ultra-quenching method, in which molten metal is supplied onto a rotating roll and solidified to produce an alloy in the form of thin strips or flakes, or fine particles such as atomization method. Segregation during solidification can be suppressed by adopting a solidification method. Further, the obtained alloy L may be subjected to heat treatment for homogenizing the structure, or to introduction of cracks and embrittlement of the alloy by hydrogen absorption/desorption.

<Step C> Step of Obtaining Mixed Fine Powder of Alloy M and Alloy L The alloy M obtained in the step A and the alloy L obtained in the step B are pulverized and mixed to obtain a mixed fine powder. In addition, when the alloy L is produced as fine particles by the atomization method, the alloy L does not necessarily need to be pulverized. The order of pulverization and mixing can be either first. The mixed fine powder may be obtained by fine pulverization after mixing at the alloying stage, or the mixed fine powder may be obtained by finely pulverizing the alloy M and the alloy L and then mixing them. The mixing ratio at the time of mixing is such that the mixing ratio of alloy L is 1.5% or more and 10% or less by weight. If the mixing ratio of alloy L is less than 1.5%, the effect of alloy L may not be sufficiently obtained, and a large amount of bcc-(Fe, Co, Ti) phase may be generated during sintering. Moreover, if the mixing ratio of the alloy L is higher than 10%, the 1-12 phase ratio in the sintered body is lowered, resulting in deterioration of the magnetic properties, which is not preferable. Also, as a method for pulverizing the alloy, a known method such as a jet mill, stamp mill, or ball mill can be adopted. Grinding is performed in an inert gas such as nitrogen, argon, or helium to suppress oxidation of the powder and reduce the risk of ignition or explosion. A small amount of air or oxygen may be mixed with the inert gas in order to improve the handleability of the fine powder after pulverization. Considering the handling and moldability of the powder, the particle size of the mixed fine powder is such that the D50 (volume-based median diameter of the particles when the cumulative frequency is 50%) obtained by the laser diffraction method by the air dispersion method is 1 μm or more. It is preferable to make it 20 μm or less. If D50 is less than 1 μm, the risk of ignition increases and the mold is damaged during molding, which is not preferable. Further, when D50 is larger than 20 μm, densification is difficult to proceed in the sintering process, which is not preferable.

<Step D> Step of Producing Green Compact The mixed fine powder obtained in Step D is molded to obtain a green compact. In order to orient the crystals, molding may be performed while applying a magnetic field during molding.

<Step E> Sintering Step A sintered body is obtained by heat-treating the green compact obtained in Step E. As a sintering method, a method in which solid phase sintering or liquid phase sintering proceeds by holding at a high temperature in a vacuum or an inert gas atmosphere, or a pressure method such as hot pressing or hot isostatic pressing (HIP). A known method including a method of holding the powder at a high temperature while applying pressure can be employed. In order to prevent oxidation due to the atmosphere during sintering, sintering is preferably performed in a vacuum atmosphere or in an inert gas such as argon or helium. Furthermore, since Sm evaporates particularly significantly at high temperatures, it is more preferable to suppress the evaporation of Sm by a method such as covering the powder compact, sealing it, or sealing it with a substance containing Sm. The sintering temperature is 900° C. or higher and 1250° C. or lower. If the sintering temperature is less than 900° C., the liquid phase will not be sufficiently formed, making it difficult to densify. Also, if the sintering temperature exceeds 1250° C., the 1-12 phase may decompose. The sintering temperature is more preferably 1000° C. or higher and 1150° C. or lower. The sintering treatment time is 5 minutes or more and 50 hours or less. If the sintering treatment time is less than 5 minutes, the densification may not proceed sufficiently. Moreover, if the sintering treatment time exceeds 50 hours, the lead time becomes long, which is not preferable in terms of operation. The pressure during pressure sintering is desirably 1000 MPa or less. Further, after the sintering process, heat treatment or diffusion treatment may be additionally performed for the purpose of improving magnetic properties.

As long as the final composition of the sintered body is set to the composition formula (1), the sintered body can be produced from only one type of alloy without using two types of alloys M and L. However, if the alloy composition is such that the final composition of the sintered body is the composition formula (1) with one type of alloy, the R-rich phase such as the 2-17 phase is larger than the 1-12 phase at the alloying stage. generate it. In that case, the ratio of the 1-12 phase is low, and the 1-12 phase is not sufficiently coarse with respect to the pulverization size, so that it is not sufficiently oriented when compacted in a magnetic field. Therefore, in order to ensure sufficient orientation in the magnetic field, alloy M and alloy L must be produced separately and mixed in the middle.

<Experimental example>
Examples of the present invention will be specifically described below, but the present invention is not limited to these examples.

(Step of producing alloy M)
First, a raw material alloy M was produced by a strip casting method. Y, Sm, Fe, Co, Ti, and Cu with a purity of 99.9% or more are used as raw material metals, and the evaporation of rare earth elements during melting is taken into account, and the final alloy composition obtained is the composition shown in Table 1. The target value was determined and weighed as follows. The weighed metals were mixed and charged into a silica crucible, and heated to 1500° C. by high-frequency induction heating to melt the raw materials. Thereafter, the molten metal was cooled to 1450° C., temporarily stored in a tundish, and then fed onto a cooling roll made of copper rotating at a peripheral speed of 1.5 m/s for cooling. The cooled alloy was pulverized by a pulverizer installed below the cooling rolls.

A part of the raw material alloy M produced was pulverized in an Ar flow chamber using a mortar and classified using a 425 μm mesh and a 75 μm mesh. Using pulverized powder with a particle size of 75 to 425 μm, component analysis of Y, Sm, Fe, Co, Ti, and Cu was performed by ICP (inductively coupled plasma) emission spectrometry. In addition, using pulverized powder having a particle size of 425 μm or more, the oxygen content was analyzed by the inert gas fusion/heat conduction method.

500 g of each of the raw material alloys M thus prepared was weighed and placed in a molybdenum container, and the container was inserted into a tubular heat treatment furnace. After replacing the atmosphere in the furnace with Ar, heat treatment was performed at 1100° C. for 1.5 hours in an atmosphere in which 2 L/min of Ar was allowed to flow. After the heat treatment was completed, the heat treatment furnace was opened to cool the alloy. At this time, the average cooling rate from 1100°C to 100°C was 10°C/min or more.

The heat-treated alloy M obtained in the above steps was embedded in a resin, polished, and the cross section of the alloy was observed with a scanning electron microscope (SEM). A JCM-6000Plus NeoScope (registered trademark) manufactured by JEOL Ltd. was used as the SEM, and BSE images were taken at an acceleration voltage of 15 kV. Using SEM image analysis software Scandium 5.2, the area ratio of each phase was calculated from the BSE image of a region of 230 μm×150 μm in the photographed image, and the average value of 2 fields of view was calculated for each alloy.

Table 1 shows the compositions of alloys M1 to M8, the values of x ₂ , y ₂ , z ₂ , w ₂ , α ₂ , β ₂ and the area ratio of 2-17 phase in the alloys after heat treatment. Alloys that satisfy the compositional formula (2) are marked with ◯, and alloys that do not satisfy the formula (2) are marked with x.

Cross-sectional BSE images of alloys M1 to M8 after heat treatment are shown in FIG. The alloys of M1 to M3 had a low 2-17 phase ratio after heat treatment, and the grain size of the 1-12 phase generally exceeded 10 μm, resulting in a coarse structure. On the other hand, in the alloys M4 to M8, the 2-17 phase was present at 10% or more even after the heat treatment, so the 1-12 phase ratio was low and the grain size of the 1-12 phase was fine.

(Step of producing alloy L)
Next, a raw material alloy L was produced. Raw material alloy no. Regarding L1, the target value is such that the final alloy composition obtained by taking into consideration the evaporation of rare earth elements during melting of raw metals of Y and Sm with a purity of 99.9% or more is as shown in Table 2. was determined and weighed. The weighed metals were mixed and put into a quartz tube, and the temperature was raised by high-frequency induction heating to melt the raw material. After that, it was cooled in a quartz tube to obtain a raw material alloy. The raw material alloy L1 thus obtained was cut in a glove box in an Ar stream atmosphere, and shavings were collected. The resulting shavings were pulverized in a mortar, the pulverized powder was sieved through 425 μm mesh and 75 μm mesh, and the powder that passed through the 75 μm mesh was collected for mixing. Moreover, raw material alloy No. For L2, Y, Sm, and Cu raw material metals with a purity of 99.9% or more were weighed so that the resulting alloy composition would be the target value, taking into consideration the evaporation of rare earth elements during melting. The weighed metals were mixed and put into a quartz tube, and the temperature was raised by high-frequency induction heating to melt the raw material. After that, the molten metal was tapped on a cooling roll made of copper rotating at a peripheral speed of 20 m/s and cooled to prepare a quenched alloy. The obtained raw material alloy L2 was pulverized using a mortar in a glove box in an Ar atmosphere, the pulverized powder was sieved with a 425 μm mesh and a 75 μm mesh, and the powder that passed through the 75 μm mesh was recovered for mixing. Of the pulverized powder of the raw material alloy L, pulverized powder having a particle size of 75 to 425 μm was used to analyze the components of Y, Sm, and Cu by ICP (inductively coupled plasma) emission spectrometry. In addition, using pulverized powder having a particle size of 425 μm or more, the oxygen content was analyzed by the inert gas fusion/heat conduction method. The compositions of alloys L1 and L2 and the values of x ₃ , w ₃ , α ₃ and β ₃ are shown in Table 2. Both alloys L1 and L2 satisfied composition formula (3).

(Step of obtaining mixed fine powder of alloy M and alloy L)
After the heat treatment, the alloy M was pulverized using a mortar in a glove box in an Ar atmosphere. The pulverized powder was sieved through a 1 mm mesh, and the powder that passed through the mesh was recovered. Zinc stearate was added to the collected pulverized powder and mixed with a rocking mixer for 15 minutes. At this time, zinc stearate was added so that the weight ratio of ground powder and zinc stearate was 100:0.035. After the heat treatment, the pulverized powder of the alloy M was pulverized using an air jet mill PJM-100 manufactured by Nippon Pneumatic Industry Co., Ltd. to obtain a fine powder. Nitrogen gas was used as the pulverizing gas, and pulverization was performed at a pulverizing pressure of 7.5 MPa. The fine powder of the raw material alloy M and the pulverized powder of the raw material alloy L which passed through a 75 μm mesh were mixed. Mixing was performed in a glove box with a nitrogen stream atmosphere, and the mixing ratios shown in Table 3 were mixed using a miller. Further, the particle size of the mixed fine powder was measured by a laser diffraction method based on an air dispersion method, and D50 of all the mixed fine powders was in the range of 5 µm or more and 6 µm or less.

(Step of producing green compact)
The mixed fine powder obtained in the above step was molded in a glove box in a nitrogen stream atmosphere. A hand press was used for molding, and a cylindrical powder compact with a diameter of 16 mm and a height of 20 mm was produced. After molding, iron capsules were filled with green compacts. The material of the iron capsule was S20C, and the inner diameter was 16 mm, the height was 20 mm, and the thickness was 2 mm.

(Sintering process)
Electron beam welding was performed in a vacuum on the iron capsule filled with the powder compact, and the capsule container and the lid were welded together to seal the capsule. A HIP treatment was performed on the sealed sample. Argon was used as a pressure medium gas, and the treatment was performed at a gas pressure of 180 MPa. The temperature was 1100° C. and the holding time was 3 hours.

The sample obtained in the above step was cut with a peripheral blade cutting machine, and the HIP body in the capsule was taken out. A portion of the HIP body was pulverized in a mortar and classified using 425 μm mesh and 75 μm mesh.

Using pulverized powder with a particle size of 75 to 425 μm, the components of Y, Zr, Sm, Fe, Co, Ti, and Cu are analyzed by ICP emission spectrometry, and the carbon content is analyzed by combustion and infrared absorption. rice field. Using pulverized powder with a particle size of 425 μm or more, the oxygen content and nitrogen content were analyzed by the inert gas fusion/heat conduction method. Also, using pulverized powder with a particle size of 75 to 425 μm, the carbon content was analyzed by the combustion/infrared absorption method. From the analysis results, the values of x ₁ , y ₁ , z ₁ , w ₁ , α ₁ , β ₁ and 1-x ₁ -2z ₁ /3-0.092α ₁ -8β ₁ /15 of each sintered body are asked.

Powder X-ray diffraction was performed using pulverized powder having a particle size of less than 75 μm. As an apparatus, a Bragg-Brentano concentrated beam type wide-angle X-ray diffractometer (X-ray diffractiometer, XRD, D8 ADVANCED/TXS manufactured by Bruker AX Co., Ltd.) was used. A Cu rotating anticathode was used as the X-ray source, the applied voltage was 45 kV, the current was 360 mA, and the Kβ filter was Ni. The scanning axis was interlocked with 2θ/θ, the interval was 0.04°, the speed was 0.6 s/step, and the range of 20°≦2θ≦80° was scanned at room temperature. From the X-ray intensity profile, the maximum intensity of the peak due to the 002 reflection of the 1-12 phase is I _ThMn12 , and the maximum intensity of the peak due to the 011 reflection of the bcc-(Fe, Co, Ti) phase is I _{α-(Fe , Co, Ti)} , the maximum intensity of the peak due to the 023 reflection of the 2-17 phase is I _Th2Ni17 , and the relative X-ray intensity of the bcc-(Fe, Co, Ti) phase is I _bcc- (Fe, Co, Ti). _Ti) 2 /I _ThMn12 and the relative X-ray intensity I _Th2Ni17 /I _ThMn12 of the 2-17 phase were determined respectively.

The cut HIP body was embedded in resin, polished, and the cross section of the HIP body was observed with a scanning electron microscope (SEM), and a local composition analysis was performed by EDX. A JCM-6000Plus NeoScope (registered trademark) manufactured by JEOL Ltd. was used as the SEM, and a BSE image was obtained at an accelerating voltage of 15 kV and EDX analysis was performed.

_The composition of _{each HIP body} produced, the types of alloys M and L used, and the mixing _ratio of alloy L are shown in Table ₃ . ₁ −2z ₁ /3−0.092α ₁ −8β ₁ /15, the values of I _{bcc-(Fe,Co,Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 , and the sintered body, the alloy M used , and alloy L satisfy the compositional formulas (1), (2) and (3) respectively.

No. 1 to 4 are experimental examples in which the mixing ratio of L1 was changed when alloy M was fixed to M1 and alloy L was fixed to L1. No. The powder XRD patterns of samples 1-4 are shown in FIG. Also, No. Values of I _{bcc-(Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 and 1-x ₁ -2z ₁ /3-0.092α ₁ - determined from the powder XRD patterns of samples 1 to 11 The relationship between the values of 8β ₁ /15 is shown in FIGS. 3 and 4. FIG. Alloy M1 has a low 2-17 phase ratio after the alloy heat treatment, and the 1-12 phase is sufficiently grown coarsely. No. L1 not mixed. No. 1 and L1 mixed at 1.1%. In No. 2, a large amount of bcc-(Fe, Co, Ti) phase was present, resulting in a high value of I _{bcc-(Fe, Co, Ti)} /I _ThMn12 . 1-x ₁ -2z ₁ /3-0.092α ₁ -8β ₁ /15 is in the range of 0 or more and 0.05 or less. Sample No. 3 has low values of I _{bcc-(Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 because the formation of bcc-(Fe, Co, Ti) phase and 2-17 phase is suppressed. result. No. 1-x ₁ -2z ₁ /3-0.092 α ₁ -8β ₁ /15 value is greater than 0.05. Sample No. 4 had a large amount of 2-17 phase and resulted in a high value of I _Th2Ni17 /I _ThMn12 . Above, No. 3 has a low 2-17 phase ratio after the alloy heat treatment, the 1-12 phase is sufficiently grown coarsely and is suitable for anisotropy, and the sintered body also has a bcc-(Fe, Co, Ti) phase. and the 2-17 phase ratio was low.

No. No. 5 to No. 7 are obtained by changing alloy M to M2 to M4, alloy L to L2, and changing the mixing ratio of L2. This is an experimental example aiming at a composition equivalent to that of Nos. 2 to 4. Alloys M2 and M3 have a low 2-17 phase ratio after the alloy heat treatment, and the 1-12 phase grows sufficiently coarsely. As a result, there are many. No. No. 2 aimed at the same composition as No. 2. 5 is No. Similar to 2, a large amount of bcc-(Fe, Co, Ti) phase was present, resulting in a high value of I _{bcc-(Fe, Co, Ti)} /I _ThMn12 . No. No. 3 aimed at the same composition as No. 3. 6 is No. Similar to 3, the values of Ibcc- _(Fe,Co,Ti) /I _ThMn12 and I _Th2Ni17 /I _ThMn12 are low because the formation of bcc-(Fe,Co,Ti) and 2-17 phases was suppressed. result. No. No. 4 aimed at the same composition as No. 4. 7 is No. Similar to 4, a large amount of 2-17 phase was present, resulting in a high value of I _Th2Ni17 /I _ThMn12 . Above, No. 6 has a low 2-17 phase ratio after the alloy heat treatment, the 1-12 phase is sufficiently grown coarsely and is suitable for anisotropy, and the sintered body also has a bcc-(Fe, Co, Ti) phase. and the 2-17 phase ratio was low.

No. 8 to 11 are samples in which the alloy M is changed to M5 to M8, the alloy L is not mixed, and the Y amount and Sm amount of the sintered body are changed. In all of the alloys M5 to M8, the 1-12 phase was finer than the fine powder size after the alloy heat treatment, and the 2-17 phase was abundant. No. 1-x ₁ -2z ₁ /3-0.092α ₁ -8β ₁ /15 is smaller than zero. Sample No. 8 contained a large amount of bcc-(Fe, Co, Ti) phase and resulted in a high value of I _{bcc-(Fe, Co, Ti)} /I _ThMn12 . No. 1-x ₃ -2z ₃ /3-0.092α ₃ -8β ₃ /15 in the range of 0 to 0.05. In samples 9 and 10, the formation of the bcc-(Fe, Co, Ti) phase and the 2-17 phase was suppressed, so the values of I _{bcc-(Fe, Co, Ti)} /I _ThMn12 and I _Th2Ni17 /I _ThMn12 gave a low result. No. 1-x ₁ -2z ₁ /3-0.092 α ₁ -8β ₁ /15 value is greater than 0.05. In 11 samples, a large amount of 2-17 phase was present, resulting in high values of I _Th2Ni17 /I _ThMn12 . As described above, when no mixing was performed, the 2-17 phase ratio was low after the alloy heat treatment, and the 1-12 phase was sufficiently grown coarsely to form a structure suitable for anisotropy. It was not possible to achieve both the bcc-(Fe, Co, Ti) phase and the low 2-17 phase ratio.

No. BSE images of cross sections of samples 1 to 11 are shown in FIG. No. A large amount of bcc-(Fe, Co, Ti) phase was observed in samples 1, 2, 5 and 8. Also, No. A large amount of 2-17 phase was observed in samples 4, 7 and 11. No. In samples 3, 6, 9 and 10, the formation of the bcc-(Fe, Co, Ti) phase and the 2-17 phase was suppressed, and the results corresponded well with the powder XRD results.

The sintered body for rare earth magnets of the present disclosure can be used for rare earth magnets.

Claims (3)

The overall composition is represented by the following compositional formula (1),
R1 _1-x1 R2 _x1 (Fe _1-y1 Co _y1 ) _w1-z1 Ti _z1 Cu _α1 O _β1 (1)
R1 is Y or Y and Gd, Y is 50 mol% or more of the total of R1,
R2 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, always contains Sm, and Sm is 50 mol% or more of the total of R2,
x ₁ , y ₁ , z ₁ , w ₁ , α ₁ , and β ₁ are each
0.3≦x ₁ ≦0.9,
0≦y ₁ ≦0.4,
0.38≦z ₁ ≦0.70,
7≦w ₁ ≦12,
0≦α ₁ ≦0.70,
0.02≦β ₁ ≦0.5, and 0≦1−x ₁ −2z ₁ /3−0.092α ₁ −8β ₁ /15≦0.05
A method for producing a sintered body for a rare earth magnet having a phase having a ThMn type ₁₂ crystal structure as a main phase, satisfying
The raw material melt is cooled, and the overall composition is represented by the following compositional formula (2),
R3 _1-x2 R4 _x2 (Fe _1-y2 Co _y2 ) _w2-z2 Ti _z2 Cu _α2 O _β2 (2)
R3 is Y or Y and Gd, Y is 50 mol% or more of the total of R3,
R4 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and must contain Sm, and Sm is 50 mol% or more of the total of R4,
x ₂ , y ₂ , z ₂ , w _{2 ,} α ₂ and β ₂ are each
0.3≦x ₂ ≦0.9,
0≦y ₂ ≦0.4,
0.38≦z ₂ ≦0.70,
_{9.4≤w2≤12.0} ,
a step of producing a rare earth magnet alloy M satisfying 0.44≦α ₂ ≦0.70 and β ₂ ≦0.5;
The raw material melt is cooled, and the overall composition is represented by the following compositional formula (3),
R5 _1-x3 R6 _x3 T3 _w3 Cu _α3 O _β3 (3)
R5 is Y or Y and Gd, Y is 50 mol% or more of the total of R5,
R6 is at least one selected from the group consisting of Sm, La, Ce, Nd and Pr, and must contain Sm, and Sm is 50 mol% or more of the total of R6,
T3 is at least one selected from the group consisting of Fe, Co and Ti,
x ₃ , w _{3 ,} α ₃ and β ₃ are each
0≦x ₃ <0.5,
_0≤w3≤3 ,
a step of producing a rare earth magnet alloy L satisfying 0≦α ₃ ≦2 and β ₃ ≦0.5;
a step of mixing the alloy L at a mixing ratio of 1.5% or more and 10% or less of the total weight to obtain a mixed fine powder of the alloy M and the alloy L;
A step of producing a green compact of the mixed fine powder;
A method for producing a sintered body for a rare earth magnet, comprising the step of heat-treating the green compact at 900° C. or higher and 1250° C. or lower for 5 minutes or longer and 50 hours or shorter to obtain a sintered body. In the powder X-ray diffraction pattern of the sintered body, the maximum intensity of the peak due to the 002 reflection of the phase having the ThMn ₁₂ type crystal structure is I _ThMn12 , which is due to the 011 reflection of the bcc-(Fe, Co, Ti) phase. When the maximum intensity of the peak to be _{Ibcc-(Fe, Co, Ti)} is
Ibcc- _{(Fe,Co,Ti) /IThMn12≤0.75}
The method for producing a sintered body for a rare earth magnet according to claim 1, satisfying In the powder X-ray diffraction pattern of the sintered body, the maximum intensity of the peak due to the 002 reflection of the phase having the ThMn ₁₂ type crystal structure is I _ThMn12 , and the maximum intensity of the peak due to the 023 reflection of the phase having the Th ₂ Ni ₁₇ type crystal structure When I _Th2Ni17 is the maximum intensity of the peak that
I _Th2Ni17 /I _ThMn12 ≦0.7
3. The method for producing a sintered body for a rare earth magnet according to claim 1, wherein

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